U.S. patent application number 13/580615 was filed with the patent office on 2012-12-20 for biodegradable polymeric networks and methods for manufacturing the same.
This patent application is currently assigned to Emory University. Invention is credited to Kenneth Gall, David Safranski, W. Robert Taylor, Daiana Weiss.
Application Number | 20120322895 13/580615 |
Document ID | / |
Family ID | 44507195 |
Filed Date | 2012-12-20 |
United States Patent
Application |
20120322895 |
Kind Code |
A1 |
Safranski; David ; et
al. |
December 20, 2012 |
BIODEGRADABLE POLYMERIC NETWORKS AND METHODS FOR MANUFACTURING THE
SAME
Abstract
Methods of manufacturing a three-dimensional, biodegradable,
thermoset polymeric network composition having desirable
degradation and mechanical properties, comprising a macromer
component cross-linked with a monofunctional acrylate-containing
component. The macromer component can comprise a
diacrylate-containing component polymerized with an
amine-containing component, wherein the molar ratio of the
diacrylate-containing component to the amine-containing component
is greater than or equal to 1.
Inventors: |
Safranski; David; (Atlanta,
GA) ; Gall; Kenneth; (Atlanta, GA) ; Taylor;
W. Robert; (Stone Mountain, GA) ; Weiss; Daiana;
(Atlanta, GA) |
Assignee: |
Emory University
Atlanta
GA
Georgia Tech Research Corporation
Atlanta
GA
|
Family ID: |
44507195 |
Appl. No.: |
13/580615 |
Filed: |
February 23, 2011 |
PCT Filed: |
February 23, 2011 |
PCT NO: |
PCT/US11/25950 |
371 Date: |
September 4, 2012 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
61307208 |
Feb 23, 2010 |
|
|
|
Current U.S.
Class: |
514/772.1 ;
522/144; 522/33; 525/419 |
Current CPC
Class: |
C08K 3/00 20130101; C08F
220/14 20130101; C08L 2203/02 20130101; C08F 290/062 20130101; C08F
222/1006 20130101; C08K 5/0025 20130101; C08K 5/101 20130101; C08G
63/91 20130101; C08K 5/17 20130101; C08L 2201/06 20130101; C08F
222/102 20200201; C08F 222/102 20200201; C08F 222/102 20200201;
C08F 222/102 20200201; C08F 299/024 20130101; C08F 2/48 20130101;
C08F 220/14 20130101; C08F 290/062 20130101; C08F 220/18
20130101 |
Class at
Publication: |
514/772.1 ;
525/419; 522/33; 522/144 |
International
Class: |
C08G 63/91 20060101
C08G063/91; C08J 3/28 20060101 C08J003/28; A61K 47/34 20060101
A61K047/34 |
Claims
1. A three-dimensional polymeric network composition, comprising: a
biodegradable macromer component photopolymerized with a
monofunctional acrylate-containing component, wherein the macromer
component comprises a diacrylate-containing component polymerized
with an amine-containing component, and wherein a molar ratio of
the diacrylate-containing component to the amine-containing
component in the macromer component is greater than or equal to
1:1.
2. The composition of claim 1, wherein the diacrylate-containing
component comprises one or more diacrylate compositions.
3. The composition of claim 1, wherein the diacrylate-containing
component comprises poly(ethylene glycol) diacrylate.
4. The composition of claim 1, wherein the diacrylate-containing
component comprises hexanediol diacrylate.
5. The composition of claim 1, wherein the diacrylate-containing
component comprises a mixture of poly(ethylene glycol) diacrylate
and hexanediol diacrylate.
6. The composition of claim 5, wherein a molar ratio of
poly(ethylene glycol) diacrylate to hexanediol diacrylate is less
than 1:1.
7. The composition of claim 1, wherein the molar ratio of the
diacrylate-containing component to the amine-containing component
in the macromer component is about 1.05:1 to about 1.25:1.
8. The composition of claim 1, wherein the amine-containing
component comprises 3-methoxypropylamine.
9. The composition of claim 1, wherein the monofunctional
acrylate-containing component comprises methyl methacrylate.
10. The composition of claim 1, wherein the composition further
comprises a photoiniator.
11. The composition of claim 10, wherein the photoiniator is
2-hydroxy-1-[4-(hydroxyethoxy)phenyl]-2-methyl-1-propanone.
12. The composition of claim 1, wherein the composition comprises
about 45 weight percent of the macromer component and about 55
weight percent of the monofunctional acrylate-containing
component.
13. The composition of claim 1, wherein the macromer component is
poly(ethylene glycol) diacrylate and hexanediol diacrylate
polymerized with 3-methoxypropylamine and the monofunctional
acrylate-containing component is methyl methacrylate.
14. The composition of claim 1, further comprising an active
agent.
15. The composition of claim 1, wherein the monofunctional
acrylate-containing component has a higher glass transition
temperature than the macromer component.
16. The composition of claim 15, wherein the glass transition
temperature of the monofunctional acrylate-containing component is
at least 20 degrees high than the glass transition temperature of
the macromer component.
17. A method of manufacturing a three-dimensional polymeric network
composition, the method comprising: polymerizing a
diacrylate-containing component with an amine-containing component
to form a biodegradable macromer component, wherein a molar ratio
of the diacrylate-containing component to the amine-containing
component is greater than or equal to 1:1; and photopolymerizing
the macromer component with a monofunctional acrylate-containing
component.
18. The method of claim 17, further comprising mixing two or more
diacrylate compositions to produce the diacrylate-containing
component.
19. The method of claim 17, further comprising disposing an active
agent in the macromer and monofunctional acrylate-containing
components.
20. The method of claim 17, wherein the diacrylate-containing
component comprises poly(ethylene glycol) diacrylate.
21. The method of claim 17, wherein the diacrylate-containing
component comprises hexanediol diacrylate.
22. The method of claim 17, wherein the diacrylate-containing
component comprises poly(ethylene glycol) diacrylate and hexanediol
diacrylate.
23. The method of claim 17, wherein the molar ratio of the
diacrylate-containing component to the amine-containing component
in the macromer component is about 1.05:1 to about 1.25:1.
24. The method of claim 17, wherein the composition comprises about
45 weight percent of the macromer component and about 55 weight
percent of the monofunctional acrylate-containing component.
25-28. (canceled)
Description
CROSS-REFERENCE TO RELATED APPLICATION
[0001] This application claims the benefit of U.S. Provisional
Patent Application Ser. No. 61/307,208, filed 23 Feb. 2010, which
is incorporated herein by reference in its entirety as if fully set
forth below.
BACKGROUND
[0002] 1. Field
[0003] The various embodiments of the present invention relate to
the syntheses and use of biodegradable polymer-based networks,
wherein the chemical composition and thermal properties of the
network may be altered to achieve specific degradation and
mechanical properties.
[0004] 2. Description of Related Art
[0005] Biodegradable polymers have been used in a multitude of
applications, such as tissue scaffolds, orthopedic devices, and
drug delivery devices. The ability to degrade is beneficial in
these applications because it enables therapeutic drugs to be
stored in the polymeric material and released without having to
surgically remove the drug-eluting device. Many of these polymers,
while biodegradable, demonstrate a loss in mechanical properties in
a relatively short period of time, which can interfere with the
overall performance of the polymeric material.
[0006] By way of example, poly(.beta.-amino ester) ("PBAE")
networks have gained attention as biodegradable polymers for use in
biomedical applications. However, many PBAE networks, while
biodegradable, are limited in thermo-mechanical properties and thus
fall apart relatively easily. The thermo-mechanical properties of
PBAE networks are mainly controlled by two parameters: (1) the
glass transition temperature ("Tg") and (2) the crosslinking
density. PBAE networks typically have a low Tg and high
crosslinking density, and thus are lacking in sufficient toughness
to survive implantation in biological applications. Further, many
of the PBAE networks of the prior art are thermoplastic materials
and thus not photopolymerizable.
[0007] Thus, polymeric materials that degrade over time, but
maintain and/or improve their mechanical properties over time would
be beneficial, particularly for biomedical applications.
BRIEF SUMMARY
[0008] Various embodiments of the present invention provide a
three-dimensional polymeric network composition, comprising a
biodegradable macromer component photopolymerized with a
monofunctional acrylate-containing component. The macromer
component can comprise a diacrylate-containing component
polymerized with an amine-containing component. In exemplary
embodiments, the molar ratio of the diacrylate-containing component
to the amine-containing component is greater than or equal to
1:1.
[0009] The diacrylate-containing component can comprise one or more
diacrylate compositions. For example, the diacrylate-containing
component can comprise poly(ethylene glycol) diacrylate (PEGDA). In
another example, the diacrylate-containing component can comprise
hexanediol diacrylate (HDDA). In yet another example, the
diacrylate-containing component can comprise a mixture of
poly(ethylene glycol) diacrylate and hexanediol diacrylate, wherein
a molar ratio of poly(ethylene glycol) diacrylate to hexanediol
diacrylate is less than 1:1.
[0010] In exemplary embodiments, the molar ratio of the
diacrylate-containing component to the amine-containing component
in the macromer component is about 1.05:1 to about 1.25:1.
[0011] In some embodiments, the amine-containing component can
comprise 3-methoxypropylamine (3MOPA).
[0012] In other embodiments, the monofunctional acrylate-containing
component can comprise methyl methacrylate (MMA).
[0013] Additionally, in some embodiments the composition can
further comprise a photoiniator. The photoiniator can be, for
example,
2-hydroxy-1-[4-(hydroxyethoxy)phenyl]-2-methyl-1-propanone.
[0014] The three-dimensional polymeric network composition can
comprise about 45 weight percent of the macromer component and
about 55 weight percent of the monofunctional acrylate-containing
component.
[0015] In some embodiments, the macromer component is poly(ethylene
glycol) diacrylate and hexanediol diacrylate polymerized with
3-methoxypropylamine, and the monofunctional acrylate-containing
components is methyl methacrylate.
[0016] The composition can further comprise an active agent.
[0017] Further, the monofunctional acrylate-containing component
has a higher glass transition temperature than the macromer
component. More specifically, the glass transition temperature of
the monofunctional acrylate-containing component is at least 20
degrees higher than the glass transition temperature of the
macromer component.
[0018] Alternative embodiments of the present invention provide a
method of manufacturing a three-dimensional polymeric network
composition by polymerizing a diacrylate-containing component with
an amine-containing component to form a biodegradable macromer
component, wherein a molar ratio of the diacrylate-containing
component to the amine-containing component is greater than or
equal to 1:1, and photopolymerizing the macromer component with a
monofunctional acrylate-containing component.
[0019] The method can further comprise mixing two or more
diacrylate compositions to produce the diacrylate-containing
component.
[0020] The method can further comprise disposing an active agent in
the macromer and monofunctional acrylate-containing components.
[0021] In some embodiments, the diacrylate-containing component can
comprise poly(ethylene glycol) diacrylate. In other embodiments,
the diacrylate-containing component can comprise hexanediol
diacrylate. In alternative embodiments, the diacrylate-containing
component comprises poly(ethylene glycol) diacrylate and hexanediol
diacrylate. More specifically, the molar ratio of the
diacrylate-containing component to the amine-containing component
in the macromer component can be about 1.05:1 to about 1.25:1.
[0022] In yet other embodiments, the three-dimensional polymeric
network composition can comprise about 45 weight percent of the
macromer and about 55 weight percent of the monofunctional
acrylate-containing component.
[0023] Alternative embodiments provide a method of therapeutically
treating a subject using a three-dimensional polymeric network
composition, the method comprising contacting a mixture with a
treatment location of the subject. The mixture can comprise a
biodegradable macromer component, a monofunctional
acrylate-containing component, and an active agent, wherein the
macromer component comprises a diacrylate-containing component
polymerized with an amine-containing component, and wherein a molar
ratio of the diacrylate-containing component to the
amine-containing component in the macromer component is greater
than or equal to 1:1. The method further comprises
photopolymerizing the macromer component with the monofunctional
acrylate-containing component of the mixture to form the
three-dimensional polymeric network having an active agent disposed
therein at the treatment location.
[0024] In some embodiments, the diacrylate-containing component
comprises poly(ethylene glycol) diacrylate and hexanediol
diacrylate.
[0025] The molar ratio of the diacrylate-containing component to
the amine-containing component in the macromer component can be
about 1.05:1 to about 1.25:1.
[0026] Further, the mixture can comprise about 45 weight percent of
the macromer component and about 55 weight percent of the
monofunctional acrylate-containing component.
[0027] The foregoing summarizes only a few aspects of the present
invention and is not intended to be reflective of the full scope of
the present invention as claimed. Additional features and
advantages of the present invention are set forth in the following
description, may be apparent from the description, or may be
learned by practicing the present invention. Moreover, both the
foregoing summary and following detailed description are exemplary
and explanatory and are intended to provide further explanation of
the present invention as claimed.
BRIEF DESCRIPTION OF THE FIGURES
[0028] FIGS. 1A-F graphically illustrate mechanical characteristics
of polymeric networks having varying PEGDA and HDDA number average
molecular weights and diacrylate to amine molar ratios, in
accordance with exemplary embodiments of the present invention.
[0029] FIGS. 2A-C graphically illustrate mechanical and degradation
properties of polymeric networks having varying PEGDA to HDDA molar
ratios, in accordance with exemplary embodiments of the present
invention.
[0030] FIGS. 3A-E graphically illustrate stress-strain curves of
polymeric networks of varying MMA weight percents, in accordance
with exemplary embodiments of the present invention.
[0031] FIG. 4 graphically illustrates glass transition temperature
of polymeric networks having varying diacrylate to amine molar
ratios as a function of MMA weight percent, in accordance with
exemplary embodiments of the present invention.
[0032] FIG. 5 graphically illustrates storage modulus of polymeric
networks having varying diacrylate to amine molar ratios as a
function of MMA weight percent, in accordance with exemplary
embodiments of the present invention.
[0033] FIG. 6 graphically illustrates storage modulus of polymeric
networks having varying diacrylate to amine molar ratios as a
function of MMA weight percent, in accordance with exemplary
embodiments of the present invention.
[0034] FIG. 7 graphically illustrates elastic modulus of polymeric
networks having varying diacrylate to amine molar ratios as a
function of MMA weight percent, in accordance with exemplary
embodiments of the present invention.
[0035] FIG. 8 graphically illustrates failure strain of polymeric
networks having varying diacrylate to amine molar ratios as a
function of MMA weight percent, in accordance with exemplary
embodiments of the present invention.
[0036] FIG. 9 graphically illustrates toughness of polymeric
networks having varying diacrylate to amine ratios as a function of
MMA weight percent, in accordance with exemplary embodiments of the
present invention.
[0037] FIG. 10 graphically illustrates failure strain of polymeric
networks having varying diacrylate to amine molar ratios as a
function of elastic modulus, in accordance with exemplary
embodiments of the present invention.
[0038] FIG. 11 graphically illustrates toughness of polymeric
networks having varying diacrylate to amine molar ratios as a
function of elastic modulus, in accordance with exemplary
embodiments of the present invention.
[0039] FIG. 12 graphically illustrates normalized mass loss of
polymeric networks having varying diacrylate to amine molar ratios
and MMA weight percents, in accordance with exemplary embodiments
of the present invention.
[0040] FIG. 13 graphically illustrates normalized mass loss of
polymeric networks having varying MMA weight percents, in
accordance with exemplary embodiments of the present invention.
[0041] FIG. 14 graphically illustrates mass loss of polymeric
networks having varying PEGDA to HDDA molar ratios, in accordance
with exemplary embodiments of the present invention.
[0042] FIG. 15 graphically illustrates mass loss of polymeric
networks of varying PEGDA to HDDA molar ratios, in accordance with
exemplary embodiments of the present invention.
[0043] FIG. 16 graphically illustrates stress-strain curves of
polymeric networks having varying MMA weight percents, in
accordance with exemplary embodiments of the present invention.
[0044] FIG. 17 graphically illustrates modulus of polymeric
networks having varying PEGDA to HDDA molar ratios, in accordance
with exemplary embodiments of the present invention.
[0045] FIG. 18 graphically illustrates stress-strain curves of
polymeric networks over varying periods of time, in accordance with
exemplary embodiments of the present invention.
[0046] FIG. 19 graphically illustrates elastic modulus over time of
polymeric networks having varying PEGDA to HDDA molar ratios, in
accordance with exemplary embodiments of the present invention.
[0047] FIG. 20 graphically illustrates failure strain over time of
polymeric networks having varying PEGDA to HDDA molar ratios, in
accordance with exemplary embodiments of the present invention.
[0048] FIG. 21 graphically illustrates toughness over time of
polymeric networks having varying PEGDA to HDDA molar ratios, in
accordance with exemplary embodiments of the present invention.
[0049] FIG. 22 graphically illustrates failure strain as a function
of rubbery modulus of polymeric networks, in accordance with
exemplary embodiments of the present invention.
[0050] FIG. 23 graphically illustrates glass transition temperature
of polymeric networks, in accordance with exemplary embodiments of
the present invention.
[0051] FIG. 24 graphically illustrates rubbery modulus of polymeric
networks, in accordance with exemplary embodiments of the present
invention.
[0052] FIG. 25 graphically illustrates failure strain as a function
of rubbery modulus of polymeric networks, in accordance with
exemplary embodiments of the present invention.
[0053] FIG. 26 graphically illustrates stress-strain curves of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
[0054] FIG. 27 graphically illustrates toughness as a function of
rubbery modulus of polymeric networks, in accordance with exemplary
embodiments of the present invention.
[0055] FIG. 28 graphically illustrates failure strain as a function
of glass transition temperature of polymeric networks, in
accordance with exemplary embodiments of the present invention.
[0056] FIG. 29 graphically illustrates failure strain of polymeric
networks, in accordance with exemplary embodiments of the present
invention.
[0057] FIG. 30 graphically illustrates toughness of polymeric
networks, in accordance with exemplary embodiments of the present
invention.
[0058] FIG. 31A-B graphically illustrates storage modulus and heat
flow of polymeric networks, in accordance with exemplary
embodiments of the present invention.
[0059] FIG. 32A-B graphically illustrates conversion of polymeric
networks, in accordance with exemplary embodiments of the present
invention.
[0060] FIG. 33 provides a nuclear magnetic resonance spectra of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
[0061] FIG. 34A-B graphically illustrates molecular weights of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
[0062] FIG. 35A-B graphically illustrates conversion of polymeric
networks, in accordance with exemplary embodiments of the present
invention.
[0063] FIG. 36A-B graphically illustrates the sol fraction of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
[0064] FIG. 37A-B graphically illustrates normalized mass loss and
water content of polymeric networks, in accordance with exemplary
embodiments of the present invention.
[0065] FIG. 38A-C graphically illustrates degradation profiles of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
[0066] FIG. 39 illustrates the development and degradation of
polymeric networks, in accordance with exemplary embodiments of the
present invention.
DETAILED DESCRIPTION
[0067] Referring now to the figures, wherein like reference
numerals represent like parts throughout the several views,
exemplary embodiments of the present invention will be described in
detail. Throughout this description, various components can be
identified as having specific values or parameters, however, these
items are provided as exemplary embodiments. Indeed, the exemplary
embodiments do not limit the various aspects and concepts of the
present invention as many comparable parameters, sizes, ranges,
and/or values can be implemented.
[0068] It should also be noted that, as used in the specification
and the appended claims, the singular forms "a," "an," and "the"
include plural references unless the context clearly dictates
otherwise. For example, reference to a component is intended also
to include composition of a plurality of components. References to
a composition containing "a" constituent is intended to include
other constituents in addition to the one named. Also, in
describing the preferred embodiments, terminology will be resorted
to for the sake of clarity. It is intended that each term
contemplates its broadest meaning as understood by those skilled in
the art and includes all technical equivalents which operate in a
similar manner to accomplish a similar purpose.
[0069] Values may be expressed herein as "about" or "approximately"
one particular value, this is meant to encompass the one particular
value and other values that are relatively close but not exactly
equal to the one particular value. By "comprising" or "containing"
or "including" is meant that at least the named compound, element,
particle, or method step is present in the composition or article
or method, but does not exclude the presence of other compounds,
materials, particles, method steps, even if the other such
compounds, material, particles, method steps have the same function
as what is named.
[0070] It is also to be understood that the mention of one or more
method steps does not preclude the presence of additional method
steps or intervening method steps between those steps expressly
identified. Similarly, it is also to be understood that the mention
of one or more components in a composition does not preclude the
presence of additional components than those expressly
identified.
[0071] The various embodiments of the present invention provide
photopolymerizable, biodegradable, three-dimensional, thermoset
polymeric networks and methods for manufacturing the same. The
polymeric networks generally comprise at least one high glass
transition temperature ("Tg"), mechanical-strength providing
component and at least one low Tg component. To increase the
thermo-mechanical properties of the network, the amount of the high
Tg, mechanical-strength providing component can be increased.
Contrastingly, to increase the degradation properties of the
network, the amount of the low Tg component can be increased.
Therefore, the polymeric networks can be manipulated to achieve
desired degradation and mechanical properties by manipulating the
balance of the components. The polymeric network disclosed herein
combines low Tg components having high crosslinking densities with
high Tg components having low crosslinking densities, such that, as
the low Tg components degrade away, the high Tg, low crosslinking
density components remain, thus increasing the mechanical
properties as the polymeric material degrades.
[0072] The photopolymerizable, biodegradable, three-dimensional,
thermoset polymeric networks disclosed herein comprise a macromer
component photopolymerized with a monofunctional
acrylate-containing component. The macromer component comprises a
diacrylate-containing component and an amine-containing component.
That is, the diacrylate-containing component and the
amine-containing component undergo step-growth polymerization to
form the macromer component. The macromer component is then
photopolymerized with the monofunctional acrylate-containing
component. An active agent, as defined below, can be added to the
macromer and/or monofunctional acrylate. A photoiniator can also be
added to the mixture of the monofunctional acrylate and/or macromer
to enable the photopolymerization of the polymeric network.
[0073] The macromer component comprises a diacrylate-containing
component polymerized with an amine-containing component and is
formed when the diacrylate-containing component and the
amine-containing undergo step-growth polymerization. In the
macromer component, the diacrylate to amine molar ratio in the
macromer component is greater than 1:1. In exemplary embodiments,
the diacrylate to amine molar ratio ranges from 1.05:1 to 1:25:1.
The diacrylate to amine molar ratio being greater than 1:1 allows
the polymeric network to photopolymerize. Specifically, such
diacrylate to amine ratios enable the amine-containing component to
completely or almost completely react with the
diacrylate-containing component, therefore leaving only diacrylate
endgroups. This is beneficial because amine endgroups prevent the
polymeric network from photopolymerizing. Contrastingly, diacrylate
endgroups enable photopolymerization of the polymeric network and
facilitate degradation.
[0074] The diacrylate-containing component of the macromer
component can comprise any diacrylate composition. Further, the
diacrylate-containing component can comprise one or more diacrylate
compositions. The diacrylate compositions can vary in number
average molecular weight. One skilled in the art will appreciate
that larger number average molecular weight diacrylate compositions
will create longer polymeric chains during the step-growth
polymerization process. Contrastingly, smaller number average
molecular weight diacrylate compositions will create shorter
polymeric chains during the step-growth polymerization process.
Further, diacrylate compositions comprise unique mechanical and
degradation properties. Therefore, the diacrylate-containing
component can be tailored to desired biological applications by
selecting a diacrylate composition or a mixture of diacrylate
compositions selected based on the desired number average molecular
weight, mechanical, and degradation properties.
[0075] Examples of diacrylate compositions include poly(ethylene
glycol) diacrylate ("PEGDA") and hexanediol diacrylate ("HDDA").
PEGDA has a higher degradation rate in comparison to HDDA due to
PEGDA's hydrophilic properties. Accordingly, one skilled in the art
will understand that PEGDA, HDDA, and mixtures thereof influence
the overall degradation properties of the polymeric network.
[0076] Exemplary embodiments of the diacrylate-containing component
comprise a mixture of PEGDA and HDDA. It shall be understood,
however, that while the examples described herein reference a
diacrylate component comprising a mixture of PEGDA and HDDA, other
diacrylate compositions, including diacrylate compositions solely
comprising PEGDA or solely comprising HDDA, can make up the
diacrylate-containing component. PEGDA has a higher degradation
rate than HDDA. Therefore, to increase the degradation rate of the
polymeric network, the amount of PEGDA in the diacrylate-containing
component should be increased. Conversely, to decrease the
degradation rate of the polymeric network, the amount of HDDA in
the diacrylate-containing component should be increased. FIGS. 1A-F
graphically illustrate the mechanical characteristics of polymeric
networks utilizing PEGDA and HDDA of varying number average
molecular weights. The data of FIGS. 1A-F provide those skilled in
the art with guidance in selecting a polymeric network having
specific mechanical properties.
[0077] Exemplary embodiments of the diacrylate-containing component
comprise a mixture of PEGDA and HDDA, wherein the PEGDA:HDDA molar
ratio is less than or equal to 1:1. For example, the PEGDA:HDDA
molar ratio can be 10:90 and 25:75. As graphically illustrated in
FIG. 2A, the fastest degradation rate of the polymeric network
occurs when the PEGDA:HDDA molar ratio is 25:75 and the slowest
degradation rate of the polymeric network occurs when the
PEGDA:HDDA molar ratio is 0:100 (i.e., solely HDDA).
[0078] As FIG. 2B graphically illustrates, the glass transition
temperatures of the polymeric networks are the highest when the
PEGDA:HDDA molar ratio is 25:75. One skilled in the art will
understand that mechanical and degradation properties are
interrelated in that one property can be manipulated at the expense
of the other property. For example, PEGDA provides mechanical
strength to the polymeric network, as illustrated in FIG. 2B, but
also increases the degradation rate of the polymeric network, as
illustrated in FIG. 2A. Therefore, if PEGDA is not included in the
diacrylate-containing component, mechanical properties are traded
in the interest of the polymeric network having a slower
degradation rate. One skilled in the art will understand that the
PEGDA:HDDA molar ratio can be manipulated to achieve desired
degradation and mechanical properties. FIG. 2C graphically
illustrates the toughness of polymeric networks having varying
PEGDA:HDDA ratios.
[0079] The diacrylate-containing component can be reacted and
polymerized with an amine-containing component to produce the
macromer component. The amine-containing component can comprise any
primary or secondary amine. For example, the amine-containing
component can comprise primary amines, methylamine, ethylamine,
butylamine, propylamine, and/or isoproplyamine. As another example,
the amine-containing component can comprise secondary amines,
dimethylamine, ethylpropylamine, and/or diethylamine. In exemplary
embodiments, the amine-containing component comprises primary
amine, 3-methoxypropylamine ("3MOPA"). While the examples described
herein specifically refer to 3MOPA, it shall be understood that any
primary or secondary amine, or combinations thereof, can be used in
the amine-containing component.
[0080] Once selected, the diacrylate-containing component and
amine-containing component can be combined to form the macromer
component. Specifically, the diacrylate-containing component and
the amine-containing component undergo a step-growth polymerization
to form the macromer component. As stated above, the
diacrylate:amine ratio of the macromer component is important
because it determines whether the polymeric network can be
photopolymerized.
[0081] The macromer component, which comprises the
diacrylate-component and the amine-containing component, can then
be photopolymerized with a monofunctional acrylate. The macromer
component generally comprises biodegradable, low Tg components, and
thus is lacking in mechanical strength. It is therefore desirable
to photopolymerize the macromer component with a monofunctional
acrylate, which improves the mechanical strength of the polymeric
network due to their high Tg properties and low crosslinking
densities. The balance between the macromer component and the
monofunctional acrylate provide a polymeric network that stiffens
as it degrades. Specifically, the balance between the macromer
component and the monofunctional acrylate enable the polymeric
network to have initial elastomeric mechanical properties that
transition to brittle and/or elastic-plastic properties as it
degrades. More specifically, the balance between the macromer
component and the monofunctional acrylate enable the polymeric
network to increase in toughness, modulus, strength, and
strain-to-failure as it degrades.
[0082] Examples of monofunctional acrylates include methyl
methacrylate, 1-hexadecyl methacrylate, 2-ethylhexyl acrylate,
2-methoxyethyl methacrylate, 2-naphthyl acrylate, 2-phenylethyl
acrylate, and 4-chlorophenyl acrylate. In exemplary embodiments,
the monofunctional acrylate is methyl methacrylate ("MMA"). It
shall be understood, however, that while MMA is described in the
various examples described herein, other monofunctional acrylates
can also be used for the monofunctional acrylate-containing
component.
[0083] Because of the unique mechanical and degradation properties
of the macromer component and the monofunctional acrylate, one
skilled in the art will understand that the weight percents of the
macromer component and the monofunctional acrylate in the polymeric
network can be manipulated to achieve desired degradation and
mechanical properties. In exemplary embodiments, the polymeric
network comprises about 55 weight percent of the monofunctional
acrylate and about 45 weight percent of the macromer component. In
other embodiments, however, the polymeric network can comprise 0-75
weight percent of the monofunctional acrylate. One skilled in the
art will understand that an increase in monofunctional acrylate
weight percent (and thus a decrease in macromer component weight
percent) will improve the mechanical properties of the polymeric
network and increase the Tg of the overall polymeric network.
Conversely, decreasing the monofunctional acrylate weight percent
(therefore increasing the macromer component weight percent) will
improve the degradation properties of the polymeric network. FIGS.
3A-E provide stress-strain curves of polymeric networks having
varying weight percents of MMA, in accordance with exemplary
embodiments of the present invention.
[0084] Once the macromer component and the monofunctional acrylate
are provided, a photoinitiator can be added to the monofunctional
acrylate and/or the macromer component. The photoiniator can be any
photoiniator, however in exemplary embodiments the photoiniator is
2-hydroxy-1-[4-(hydroxyethoxy)phenyl]-2-methyl-1-propanone.
[0085] An "active agent" can also be included in the polymeric
network. As used herein, "active agent" means a pharmaceutical or
biotechnological compound or construct that induces a biological,
pharmacological, or cosmetic effect on an organism. An active agent
can be a compound, molecule, chemical, or biological construct that
provides a physical or chemical change to an existing
condition.
[0086] The polymeric network of the present invention permits the
delivery of active agents, including therapeutics, diagnostics, and
prophylactics that may or may not be delivered using polymeric
networks currently known in the art. Active agents of the present
invention include, but are not limited to: agents for gene therapy;
nucleic acids; DNA; RNA; polynucleotides; peptides; proteins; amino
acids; carbohydrates; viruses; antigens; immunogens; antibodies;
chemical or biological materials or compounds that induce a desired
biological or pharmacological effect; anti-infectives, such as
antibiotics and antiviral agents; analgesics and analgesic
combinations; anorexics; antihelminthics; antiarthritics;
antiasthmatic agents; anticonvulsants; antidepressants;
antidiabetic agents; antidiarrheals; antihistamines;
antiinflammatory agents; antimigraine preparations; antinauseants;
antineoplastics; antiparkinsonism drugs; antipruritics;
antipsychotics; antipyretics; antispasmodics; anticholinergics;
sympathomimetics; xanthine derivatives; cardiovascular preparations
including potassium and calcium channel blockers, beta-blockers,
alpha-blockers, and antiarrhythmics; antihypertensives; diuretics
and antidiuretics; vasodilators including general coronary,
peripheral and cerebral; central nervous system stimulants;
vasoconstrictors; cough and cold preparations, including
decongestants; hormones, such as estradiol and other steroids,
progesterone and derivatives, testosterone and derivatives;
corticosteroids; angiogenic agents; antiangeogenic agents;
hypnotics; immunosuppressives; muscle relaxants;
parasympatholytics; nicotine; psychostimulants; sedatives;
tranquilizers; ionized and nonionized active agents; cells; and
compounds of either high or low molecular weight, among others. An
active agent can further comprise a particle or plurality of
particles, wherein a particle may induce a biological,
pharmacological or cosmetic effect on an organism. Particles can
comprise metals, non-metals, ceramics, polymers, organics,
inorganics, composites, or combinations thereof. Examples of
particles comprise, but are not limited to, liposomes, viruses,
polymer particles that encapsulate active agents, which are
released over time, coated particles that facilitate delivery of an
active agent, wherein the particles comprise gold, polystyrene,
glass, tungsten, platinum, ferrite, glass, or latex, among others.
The active agents may have local effects, such as providing for a
local anesthesia, or may have systemic effects.
[0087] One skilled in the art will understand that active agents
can be of various weight percents in the polymeric network. For
example, in some embodiments, the active agent can comprise up to
70 weight percent of the polymeric network. It shall also be
understood that degradation and mechanical properties of the
polymeric network directly influence how the active agent is
delivered to aid in the treatment of a subject. For example, the
polymeric network can be designed to provide a quick, anesthetic
delivery of the active agent to the patient or to provide slow,
long-term delivery of the active agent to the patient.
[0088] The polymeric network can then be photopolymerized. Because
the polymeric network is a thermoset material, it can be injected
into or disposed on a treatment site of a subject and cured in
vivo. The curing of the polymer network disclosed herein do not
release the high heats of those in the prior art, thus providing
advantages to doctors and the like who want to directly inject and
cure the polymeric network in vivo. The development and degradation
of the polymeric networks is illustrated in FIG. 39.
EXAMPLES
[0089] The various embodiments of the present invention are
illustrated by the following non-limiting examples.
Example 1
Polymeric Networks Comprising HDDA, 3 MOPA, and MMA
[0090] Acrylate-terminated macromers were formed via a step-growth
polymerization of HDDA and 3MOPA at molar ratios of 1.05:1 to
1.20:1. The step-growth polymerization reaction proceeded at about
90.degree. C. for about 24 hours on a rotary shaker at about 200
rotations per minute (rpm). The photoiniator, Irgacure 2959, was
added at 0.5 weight percent and the polymer network were
photopolymerized with a UV lamp at 365 nanometers (nm). MMA was
added to select macromers at varying weight percent ratios prior to
photopolymerization.
[0091] Materials of 1 square centimeter (cm.sup.2) cut from 1
millimeter (mm) thick sheets were degraded in phosphate buffered
saline, pH=7.0, at 37.degree. C. for 12 weeks. Samples were massed
at predetermined times and removed from the phosphate buffered
saline. Samples were then dried for 24 hours and massed again to
determine the mass loss.
[0092] Two methods were used for mechanical characterization.
Dynamic mechanical analysis (TA Q800) determined the
thermomechanical properties from about -100.degree. C. to
100.degree. C. at a rate of about 3.degree. C./minute. The same
were run in tension under strain control of about 0.1%. Standard
tensile testing (MTS Insight 2) of ASTM type IV dogbone samples
occurred at a strain rate of 10.sup.-3 at about 37.degree. C. in a
Thermcraft chamber to determined the bulk mechanical
properties.
[0093] Since the macromer base without MMA is highly crosslinked,
it has rather poor mechanical properties at the temperature of
interest. Further, the Tg is far below the operating temperature,
thus the mechanical properties are rather week. The initial Tg is
near -40.degree. C., thus the Tg needed to be increased some
70.degree. C. to be near body temperature. The glass transition
temperature was increased by adding MMA, as shown in FIG. 4.
[0094] By varying the diacrylate to amine molar ratio from 1.05:1
to 1.20:1, the molecular weight of the resultant macromer can be
varied. Four different molecular weight macromers were mixed with
MMA. The modulus of these materials can be found in FIG. 5. The
modulus starts to greatly increase beyond 45 weight percent (wt %)
MMA. The materials remain rubbery below 45 wt % MMA, then
transition to a glassy rigid plastic beyond 45 wt % MMA.
[0095] If the modulus is taken at a relative temperature from the
Tg and in the rubbery regime, the curves become normalized. The
storage modulus at a temperature of Tg+75.degree. C. for each
composition as a function of weight percent MMA is shown in FIG. 6.
It is expected that there will be a decrease in the rubbery modulus
because adding MMA to the network decreases the crosslinking
density, therefore decreasing the rubbery modulus.
[0096] FIGS. 7-11 were produced from stress-strain curves via
Insight 2 tensile tester. FIG. 7 shows the increase in elastic
modulus as wt % of MMA is increased for four molecular weight
macromers. The elastic modulus did not change until near 55 to 60
wt % MMA. While this varies slightly from FIG. 5, the elastic
modulus of a tensile tester varies slightly from the modulus of a
DMA, the variation is often less than an order of magnitude. FIG. 8
shows the relationship between failure strain and MMA
concentration. The failure strain increases as the MMA
concentration increases because the Tg is increasing. FIG. 9 shows
the relationship between toughness and MMA concentration. Again,
toughness increases as MMA concentration increases because the
failure strain is increasing. FIGS. 10 and 11 show the
relationships between (1) failure strain and modulus and (2)
toughness and modulus, respectively.
[0097] In FIG. 12, the degradation profile of four macromers at a
set wt % MMA allowed the determination of the effect of molecular
weight of the macromer on degradation rate. There are no
considerable differences between the four networks at a set wt %
MMA, thus wt % MMA is the determining factor in degradation. In
FIG. 13, the degradation profile comprises one macromer at three
different wt % MMA. As the wt % MMA increases, the degradation rate
and degree of degradation over 12 weeks decreases. Thus the wt %
MMA is a major component in controlling the rate of degradation for
these networks.
Example 2
Polymeric Networks Comprising PEGDA, Hdda, 3 Mopa, and MMA
[0098] With the addition of PEGDA, M.sub.n.about.700, to HDDA
during the step-growth polymerization reaction, a copolymer was
formed. By varying the molar ratio of PEGDA:HDDA, the degradation
rate and water uptake can be tuned. The tuning of the degradation
rate is illustrated in FIG. 14. The addition of MMA allows for a
suppression of the increased degradation rate, as illustrated in
FIG. 15.
[0099] Exemplary stress-strain curves of HDDA-MMA networks are
shown in FIG. 16. The failure strain, strength, and toughness
increases as MMA concentration is increased, as described above. It
is expected that over time, the networks would degrade, and their
mechanical properties would decrease, as seen in networks without
MMA in FIG. 17. Tensile samples were degraded in saline at about
37.degree. C. in an incubator for 8 weeks. Exemplary stress-strain
curves are shown in FIG. 18. The strengthening is due to a relative
shift in Tg. As time increases, the network degrades, losing its
low Tg network structures. The remainder of the network has a
higher relative Tg, thus showing the transition from a more rubbery
material to a visco-elastic material. FIGS. 19-21 detail the
mechanical properties as a function of time for DDA-PEGDA-MMA
networks. As PEGDA concentration increases, the network degrades
faster, and has a shift to a higher Tg and increased mechanical
properties. Thus, the addition of PEGDA to the MMA network allows
further tailorability of mechanical properties during degradation.
This is significant because it demonstrates a polymer that
self-toughens as it degrades.
Example 3
Effect of Chemical Structure and Crosslinking Density on the
Thermo-Mechanical Properties and Toughness of (Meth)Acrylate Shape
Memory Polymer Networks
[0100] Sixteen mono-functional (meth)acrylates were used as linear
chain builders and 16 multi-functional (meth)acrylates were used as
the crosslinkers to form the polymer networks. The names,
abbreviations, chemical structures, and molecular weights can be
found in Charts 1 and 2. A set of networks comprised of 10 mole
percent (mol %) PEGDMA550 were copolymerized with each
monofunctional acrylate from Chart 1. A set of networks comprised
of 10 mol % of each crosslinker from Chart 2 were copolymerized
with 90 mol % tBA. These sets were calculated using the molecular
weights given in Charts 1 and 2. In addition, equivalent molar
amounts of BMA, tBA, and EEM were copolymerized in varying degrees
with PEGDMA550. The photoinitiator,
2,2-dimethoxy-2-phenylacetophenone, was added to each material in
an amount of 0.5 wt %. Further equivalent molar amounts of BZA and
EGPEM were copolymerized with PEGDMA550. All materials were
purchased from Sigma Aldrich or Polysciences and used as
received.
TABLE-US-00001 CHART 1 Mono-functional Monomers Monomer Structure
Molecular weight (g/mol) Methyl acrylate (MA) ##STR00001## 86.09
Methyl methacrylate(MMA) ##STR00002## 100.12 Butyl acrylate(BA)
##STR00003## 128.17 tert-Butyl acrylate(tBA) ##STR00004## 128.17
tert-Butyl methacrylate(tBMA) ##STR00005## 142.20 2-Ethoxyethyl
methacrylate (EEM) ##STR00006## 158.19 Isobornyl methacrylate (IMA)
##STR00007## 222.32 2-Ethylhexyl methacrylate (2EHM) ##STR00008##
198.3 Isodecyl acrylate (IA) ##STR00009## 212.33 Benzyl
methacrylate (BMA) ##STR00010## 176.21 Ethylene glycol phenyl ether
methacrylate (EGPEM) ##STR00011## 206.24 Poly(propylene glycol)
acrylate (PPGA) ##STR00012## 547 Poly(ethylene glycol) phenyl ether
acrylate M.sub.n 236 (PEGPEA236) ##STR00013## 236 Poly(ethylene
glycol) phenyl ether acrylate M.sub.n 280 (PEGPEA280) ##STR00014##
280 Poly(ethylene glycol) phenyl ether acrylate M.sub.n 324
(PEGPEA324) ##STR00015## 324 Benzyl acrylate (BZA) ##STR00016##
162.2
TABLE-US-00002 CHART 2 Multi-functional Monomers Molecular Monomer
Structure Weight(g/mol) Bisphenol A ethoxylate dimethacrylate
M.sub.n 1700 (BPA1700) ##STR00017## ~1700 Bisphenol A ethoxylate
dimethacrylate M.sub.n 540 (BPA540) ##STR00018## ~540 Bisphenol A
ethoxylate diacrylate M.sub.n 688 (BPA688) ##STR00019## ~688
Bisphenol A ethoxylate diacrylate M.sub.n 512 (BPA512) ##STR00020##
~512 Bisphenol A ethoxylate diacrylate M.sub.n 468 (BPA468)
##STR00021## ~468 Neopentyl glycol propoxylate diacrylate (NGPDA)
##STR00022## 328 1,6-Hexanediol diacrylate (HEXDA) ##STR00023## 226
Poly(ethylene glycol) dimethacrylate M.sub.n 550 (PEGDMA550)
##STR00024## 550 Pentaerythritol triacrylate (PETA) ##STR00025##
298 Trimethylolpropane ethoxylate triacrylate M.sub.n 428 (TETA428)
##STR00026## ~428 Trimethylolpropane ethoxylate triacrylate M.sub.n
604 (TETA604) ##STR00027## ~604 Trimethylolpropane ethoxylate
triacrylate M.sub.n 912 (TETA912) ##STR00028## ~912
Trimethylolpropane propoxylate triacrylate (TPTA) ##STR00029## ~644
Glycerol propoxylate triacrylate (GPTA) ##STR00030## ~428
Di(trimethylolpropane) tetraacrylate (DTTA) ##STR00031## 466
Dipentaerythritol penta/hexaacrylate (DPPHA) ##STR00032## 524
[0101] The polymer solutions were injected into a mold composed of
two glass slides separated by 1 millimeter (mm) spacers. Glass
slides were cleaned with Alconox then coated with Rain-X as a mold
release agent. The injected molds were polymerized under a 365
nanometers (nm) UV lamp for an average of 20 minutes (min), while
materials with low concentrations of crosslinker could take over 30
min.
[0102] Samples for dynamic mechanical analysis (DMA) were prepared
by laser cutting specimens to 20 mm.times.5 mm.times.1 mm from bulk
material. ATAQ800 was used in tensile loading with strain of 0.2%,
preload of 0.001 N, force track of 150%, and frequency of 1 Hz. The
samples were equilibrated at -50.degree. C. for 2 min then raised
to 200.degree. C. at a rate of 5.degree. C./min (n.gtoreq.2). The
glass transition temperature was defined as the peak of the tan d
curve from the DMA testing.
[0103] Mechanical tensile testing was performed on dogbones of half
size ASTM D638 type IV, which was laser cut from 1 mm thick
samples. The testing apparatus was an MTS Insight 2 mechanical
tester with a 100 N load cell. A thermal chamber (Thermcraft, Inc.,
model LBO-14-8-5.25-1X-J8249.sub.--1A) was used to isothermally
test either at the glass transition temperature of each material or
at another specified temperature. Once the chamber reached the set
temperature, 10 min were given to insure equilibrium. A
displacement rate of 1 mm/min was used, and the displacement was
measured by the crosshead. Toughness was calculated by integrating
the area under each stress-strain curve using the trapezoidal rule.
The Kendall rank correlation coefficient was calculated to describe
the relationship between select thermo-mechanical properties.
[0104] The characteristic ratios, C.infin., from Table 1 were
calculated using the method according to Wu by the following
equation:
C.sub..infin.+(1/.phi..sub.0).sup.2/3[(.SIGMA.K.sub.i+Bn.sub.r].sup.4/3(-
M.sub.v].sup.4/3(M.sub.v/l.sub.v.sup.2).
[(.SIGMA.K.sub.i+Bn.sup.r)/M.sub.r].sup.4/3 takes into account the
intrinsic viscosity of the chain, where PKi sums the molar
stiffness of each group. The molar stiffness constants for each
group such as acrylic group or phenyl rings are detailed in the
source. B takes into account the tacticity of the chain, for
example, for poly(methyl methacrylate) polymerized by free radical
polymerization, B.about.4.12. The CED for five mono-functional
(meth)acrylates was calculated using the group contribution method
outlined by Van Krevelen. The molar volume (V.sub.g) values used
were for glassy amorphous polymers. The cohesive energy was
calculated from the molar attraction values (F) using
CED=(F/V.sub.g).sup.2. Table 2 contains the calculated values. The
monomers with aromatic sidegroups had higher CED values than the
monomers with aliphatic side groups.
TABLE-US-00003 TABLE 1 Characteristic Ratios of Mono-functional
Monomers Mono-Functional Monomer C.sub..infin. tBA 9.47 EEM 11.98
BZA 12.97 BMA 13.67 EGPEM 16.19
TABLE-US-00004 TABLE 2 CED of Select Mono-functional Monomers.
Monomer CED (MPa) BMA 396 BZA 424 EGPEM 401 EEM 358 tBA 332
[0105] Materials were initially screened by creating a series of
networks with either set multi-functional crosslinker or set
monofunctional linear builder. The 16 networks in Table 3 were
produced by polymerizing 10 mol % of PEGDMA550 and 90 mol % of each
mono-functional monomer. The Tg and rubbery modulus (E.sub.r) were
measured through DMA and showed a medium strength positive
correlation. The Tg of the networks ranged from -29 to 112.degree.
C., and the E.sub.r ranged from 2.75 to 17.5 MPa. Generally, the Tg
increased as the pendant length decreased or by the addition of an
.alpha.-methyl group. The 16 networks in Table 4 were produced from
90 mol % tBA and 10 mol % of each crosslinker. The Tg and the
E.sub.r showed a medium strength positive correlation. The Tg
ranged from -2 to 98.degree. C., and the E.sub.r ranged from 6.48
to 129.5 MPa. As the functionality of the crosslinker increased,
the E.sub.r increased for equivalent mole fraction of crosslinking
molecule. The increase in rubbery modulus is driven by the relative
increase in mole fraction of crosslinking "bonds" for a crosslinker
with higher functionality.
TABLE-US-00005 TABLE 3 Thermo-Mechanical Properties of Networks
composed of 10 mol % PEGDMA550 with 90 mol % Mono-functional
(meth)acrylate. Mono-functional (meth)acrylate T.sub.g (.degree.
C.) E.sub.r (MPa) MMA 91.3 17.5 MA 23.5 11.75 BA -15 7.3 tBA 40.5
10.7 tBMA 89.5 8.9 EEM 19.5 11.25 IMA 112 6.45 2EHM 20.5 7.7 BZA 23
10.51 IA -23.5 6.5 BMA 68 9.4 EGPEM 40.5 12.75 PPGA -29 2.75
PEGPEA236 10.5 6.1 PEGPEA280 -3.5 6.05 PEGPEA324 -9.5 4.45
[0106] The 16 networks from Table 4 were tensile tested until
failure to characterize their large strain mechanical properties
including failure strain and toughness. The failure strain of each
network is plotted against its corresponding E.sub.r from DMA in
FIG. 22. The failure strain ranged from less than 10% to over a
100%. The numbers 2-5 in FIG. 22 highlight the functionality of the
crosslinkers. As expected, as the E.sub.r of the network decreases
the failure strain increases. For most crosslinkers, as the
functionality of the crosslinker decreases, the rubbery modulus
decreases, and the failure strain increases. Consistent with
previous results, a significant effect of the crosslinker chemistry
was not observed aside from property values governed by a change in
crosslinking effectiveness measured through rubbery modulus.
TABLE-US-00006 TABLE 4 Thermo-mechanical Properties of Networks
composed of 90 mol % tBA and 10 mol % Multi-functional
(meth)acrylate. Multi-functional (meth)acrylate T.sub.g (.degree.
C.) E.sub.r (MPa) BPA1700 -2.75 7.35 BPA540 70.5 8.15 BPA688 43.5
8.25 BPA512 64.5 9.0 BPA468 59.5 8.8 NGPDA 62.5 6.48 HEXDA 68.5
10.85 PEGDMA550 40.5 10.7 PETA 98 42.5 TETA428 83 25 TETA604 55
16.65 TETA912 24.5 15.95 TPTA 58 23 GPTA 69.5 15.5 DTTA 92 49.5
DPPHA 74 129.5
[0107] Five linear (meth)acrylates were selected based on their
differences in chemical structure and initial thermo-mechanical
testing data. As the crosslinker concentration was decreased, the
E.sub.r decreased. As the concentration of crosslinker approaches
zero, the E.sub.r plateau disappears and E.sub.r steadily decreases
with increasing temperature. The Tg of each network increased as
the concentration of crosslink E.sub.r decreased. A non-linear
trend is observed in FIG. 23. FIG. 24 displays the trend of the
decreasing E.sub.r as the crosslinker concentration decreased for
the five systems. Systems start at the same point since each was
originally composed of 100% PEGDMA550. Systems approach 0 MPa as
the crosslinker concentration approaches 0%. The results in FIGS.
23 and 24 demonstrate one of the known advantages of commercially
available (meth)acrylate systems; using combination of various
linear monomers and crosslinkers, one can independently tailor
glass transition temperature and rubbery modulus. It is important
to note that the PEGDMA550 crosslinker has equivalent impact on the
five selected mono-functional monomers in terms of crosslinking
effectiveness measured through rubbery modulus.
[0108] The networks were tensile tested to large strains to
understand the effect of structure on the large strain behavior of
the networks. The failure strain of each composition from the
tensile test was plotted against its respective E.sub.r from DMA in
FIG. 25. The results were plotted against E.sub.r to eliminate any
differences that may be a result of different "effective" crosslink
density in the networks and thus isolate the effects of the linear
monomer chemistry as a function of increasing crosslinker
concentration. In addition, all tests in FIG. 25 were conducted at
the Tg of the respective polymer (which differed significantly,
vis-a-vis FIG. 23) to assure all networks were in an equivalent
state of macromolecular motion. At E.sub.r greater than 10 MPa
(high crosslink density) the five systems had comparable failure
strains for all compositions. At E.sub.r lower than 10 MPa the
network failure strains diverged significantly. As the E.sub.r
further decreased below 1 MPa, the networks did not display
reliable rubbery plateaus, thus the data were excluded. The
correlation coefficients between failure strain and E.sub.r reveal
the high inverse correlation between failure strain and
E.sub.r.
[0109] To further support the results, FIGS. 25 and 26 display
representative stress-strain curves of the five systems with
increasing rubbery moduli. For all five materials, as E.sub.r
decreases, the failure strain increases. The tBA, EEM, BZA, and
EGPEM also show a decrease in strength as E.sub.r decreases. Unlike
the other systems, the BMA system does not show a steady decrease
in strength as E.sub.r decreases. The BMA has relatively higher
failure strains and failure strengths as compared to the other
materials at roughly equivalent rubbery modulus.
[0110] FIG. 27 displays the toughness, calculated as the area under
stress-strain curves of the systems, as a function of the E.sub.r.
The systems have similar toughness at relatively higher E.sub.r
values, and the systems diverge at E.sub.r values below 10 MPa. The
tBA, EEM, BZA, and EGPEM systems have toughness values nearly a
third of BMA. The point of divergence, the shape of the BMA
stress-strain curves, and the increased toughness are points of
interest to be further studied.
[0111] Networks composed of 2.5 mol % PEGDMA550-co-BMA or
PEGDMA550-co-tBA were tensile tested across a range of
temperatures, represented in FIG. 28. The objective of this testing
was to verify that the relatively high toughness of the BMA
material compared to tBA was not merely an artifact of a relative
test temperature difference even though both materials were tested
at their Tg defined as the peak in tan d. The strain to failure in
FIG. 28 is plotted at temperatures relative to each composition's
respective Tg, T-Tg. A peak in failure strain is seen 15-20.degree.
C. before the Tg, then the curves level off when well into their
respective rubbery region. The PEGDMA550-BMA curve reaches a higher
peak and is broader than the PEGDMA550-tBA curve, highlighting the
inherent toughness difference in the two materials that is not
driven by a difference in effective crosslink density (measured
through rubbery modulus) or temperature relative to Tg.
[0112] Mixtures of the various linear monomers were created with
equivalent crosslinker concentration to determine how mechanical
properties evolved from one network to another. FIG. 29 shows the
failure strain as a function of mol % BMA in three other linear
monomers (all materials contain 2.5 mol % crosslinker). As the
concentration of BMA increases, the failure strain increases. This
trend is also seen in FIG. 30, which describes the effect of
increasing the concentration of BMA on the toughness of the
networks.
[0113] Polymer networks based on (meth)acrylate monomers have
potential for a broad range of thermo-mechanical properties, making
them strong candidates for shape memory materials. In order to
understand the role of various components of these networks,
mono-functional and multi-functional (meth)acrylates were used to
synthesize a diverse set of polymer networks. Structure-property
relationships were determined in these networks by studying their
thermo-mechanical transitions and stress-strain response for
systematically varied monomer functionalities, concentrations, and
chemistries.
[0114] By holding crosslinker concentration constant, the effect of
the mono-functional (meth)acrylate structure on the networks
properties was determined. Chain backbone stiffness (capacity for
conformational motion) and cohesive energy between chains are the
primary drivers for Tg, but crosslinking and other factors also
participate. The mono-functional (meth)acrylates with long
sidegroups had the lowest Tg as may be expected based on the
reduction of steric hindrance to conformational motion from the
methylene and ester groups. As the sidegroup length decreased and
.alpha.-methyl side groups were added, the Tg increased due
primarily to local steric hindrance of segmental conformational
motion and increased cohesive energy between chains. The effects
are clear when combining the structures in Chart 1 with the Tg data
from Table 3. Since these (meth)acrylates all have the same
backbone, the sidegroup structure determines the Tg, and similar
results in epoxies have demonstrated that the chemical structure of
the amine alters Tg. In summary, the combination of both
.alpha.-methyl groups and short, rigid pendant groups on each side
of the chain's backbone increases the Tg as can be seen in MMA and
IMA.
[0115] In order to understand the effect of the crosslinker
functionality on the networks, the mono-functional acrylate, tBA,
was held constant and polymerized with various crosslinkers. The
most identifiable trend was the relationship between the
crosslinkers' functionality and Er. It is known that as the
crosslinkers' functionality increases, the network crosslink
density increases, thus increasing Er. This trend is clear in FIG.
22, where the failure strain is plotted against the Er. Driven by
different crosslinking effectiveness, the 16 networks trade-off
failure strain and rubbery modulus. The majority of the networks
with low Er had higher failure strains than the high Er networks.
The materials with high Er due to higher functionality were
relatively brittle due to high crosslink density.
[0116] The above results highlight the capacity to readily adjust
thermo-mechanical properties, a capacity that is central to an
effective shape memory polymer. Aside from basic thermo-echanical
properties, it is important for some shape memory applications, and
for deeper fundamental understanding, to examine large strain
behavior of the networks. Prior work has examined the effect of
varying crosslinker length and concentration on the large strain
behavior of acrylate networks. Here we focus on the reciprocal
problem of varying mono-functional monomer for the same crosslinker
added in varying concentrations. Five mono-functional monomers were
chosen for differences in their transition temperatures, chemical
structure, C.infin. and CED values.
[0117] In order to determine an appropriate testing temperature and
provide a rough measure of effective crosslink density, Tg and
E.sub.r were measured for all five materials across all crosslink
densities. As expected, the E.sub.r decreases as the concentration
of the crosslinker decreases in all networks. Since the selected
crosslinker (PEGDMA550) has a relatively low Tg value when
homopolymerized, the addition of it to all linear monomers serves
to reduce Tg while increasing rubbery modulus. At 1 mol %
crosslinker, the networks had reached their final Tg, thus further
characterization was not continued for the BZA and EGPEM systems.
Also, below a 1 mol % crosslinker concentration, the networks start
to effectively transition to a thermoplastic, which is indicated by
a loss of a rubbery modulus plateau. The breadth of the transition
from the glassy to rubbery state decreases as the concentration of
crosslinker decreases, as is expected because highly crosslinked
systems have increased heterogeneity. The results here are
consistent with previous studies where concentration of crosslinker
was varied in acrylates.
[0118] The baseline thermo-mechanical experiments were necessary to
assure that the selected test temperature is in the same proximity
of an individual composition's Tg and maintain equivalent states of
molecular motion during large strain testing. A key finding of the
tensile test was the existence of a divergence point, seen in FIG.
25 at a rubbery modulus of 10 MPa. Above 10 MPa, the crosslinking
dominates the large strain mechanical properties of the network and
a relatively brittle response is observed. It is important to note
that although the mono-functional monomer has minimal impact on
mechanical properties at these high crosslink densities, the
mono-functional monomer choice will influence Tg of the network and
consequently impact mechanical properties at a constant testing
temperature. As E.sub.r is decreased below 10 MPa, the large strain
mechanical properties of the networks diverge and the capacity for
strain and toughness depends on the choice of mono-functional
monomer. Soon after entering the regime of mono-functional monomer
sensitivity, the Tg of each network has reached close to a steady
state value and thus there is no correlation between the absolute
Tg of the network and the failure strain. This is evident in tBA
and EGPEM having similar Tg's at low mol % PEGDMA550, but different
failure strains at similar concentrations of PEGDMA550.
[0119] The stress-strain curves at representative rubbery moduli
values were examined to understand the divergence of the failure
strain. In general, the networks transition from brittle to ductile
behavior as the Er decreased is seen in FIG. 26. An inherent
trade-off between strength and failure strain is evident in most
networks with exception to the BMA network, which reached a high
enough strain to exhibit non-linear strain-hardening even at
reasonably high crosslink densities. This can be attributed to the
reorientation of chains in the tensile direction. As Er decreases
it becomes increasingly important to consider structural parameters
of the mono-functional monomers. The strain to failure results do
not correlate inversely with C.infin. values for the crosslinked
networks as is common for thermoplastics. For example the C.infin.
value for tBA is significantly lower than C.infin. for BMA although
the latter has significantly higher failure strain at equivalent
rubbery modulus. This observation implies that the capacity for
network backbone chains to coil, as measured by C.infin., is
incapable of predicting failure strain and toughness properties
once these chains are moderately crosslinked. It seems that factors
that toughen thermoplastics, such as coilability and high
entanglement density are rendered less effective due to chemical
crosslinking. On the other hand, the CED may be used for relative
comparison to determine if a material will strain farther through
enhanced network toughness, as seen by combining Table 2 and FIG.
25. These results indicated that higher cohesive energy between
chains, for equivalent crosslink density, serves to toughen the
materials through increased resistance to fracture during large
strain deformation. In other words, it appears that in the presence
of moderate chemical crosslinking, strain to failure can be
enhanced through improved toughness by increasing CED between
chains.
[0120] Toughness was explicitly evaluated because of its importance
during processing of shape memory polymers. Similar to failure
strain, toughness diverges at 10 MPa, as seen in FIG. 27. Due to
the strain-hardening that is observable in the stress strain
behavior, BMA has the highest toughness below the divergence point
while the other linear monomers have the same lower amount of
toughness. The parameter C.infin. also breaks down when examining
network toughness. For example, from Table 1 and FIG. 27, BZA,
EGPEM, and EEM have different calculated C.infin., but exhibit
similar levels of toughness.
[0121] In order to verify the inherently superior large strain
mechanical properties of BMA networks, the test temperature should
be eliminated as a factor influencing mechanical properties. To
assure test temperature was not a factor in comparison of the
networks, PEGDMA550-co-BMA and PEGDMA550-co-tBA, at the same mol %
crosslinker (and the same rubbery modulus), were tested over a wide
temperature range. These two materials were chosen because their
failure strains and test temperatures differed by 100% and by more
than 10.degree. C., respectively. Considering a sweep of test
temperatures, the PEGDMA550-co-BMA network has inherent capacity
for more deformation as observed in FIG. 28. It is interesting to
note that the enhanced toughness of the BMA network only occurs at
temperatures in the range of Tg -10.degree. C. to Tg+50.degree. C.
Thiol-ene/acrylate networks containing phenyl rings via Bisphenol A
ethoxylate diacrylates have shown increased impact toughness near
their Tg. In the extreme temperature limits (glassy or rubbery) the
failure strain of both materials is low and comparable. This result
indicates that the toughening mechanism has an inherent viscous
component that operates on distinct time and temperature
scales.
[0122] To ascertain the influence of varying amounts of
mono-functional monomers on mechanical properties, binary mixtures
of mono-functional monomers with constant crosslinker concentration
were formulated. With the BMA network as an upper bound of
properties, the failure strain and toughness rise as BMA
concentration increases, seen in FIGS. 29 and 30. The BMA-BZA and
BMA-EGPEM mixtures have higher failure strains and toughness values
than the BMA-tBA mixtures, which may be due to the higher and more
similar CED values of the monomers containing phenyl rings. The
mechanical properties converge as the mol % BMA increases, near 70
mol % BMA. The properties of the BMA-BZA mixtures increase as the
concentration of the .alpha.-methyl group increases, suggesting
that the increased steric hindrance from the .alpha.-methyl group
affects the mechanical properties. Likewise, the properties of the
BMA-EGPEM mixtures increase as the phenyl ring is moved closer to
the backbone by the subtraction of flexible ethylene glycol groups.
Given these two trends, the transition from tBA to BMA is
significant because both .alpha.-methyl and phenyl ring groups are
being added to the network with increased BMA concentration. tBA
lacks substantial deformation capacity because the failure strain
and toughness do not increase until the majority of the network is
BMA.
[0123] A method to theoretically predict (meth)acrylate network
properties based upon the chemistry and structure has yet to be
established. From this study, properties such as failure strain,
toughness, Tg, and E.sub.r can be tailored by varying the
components of the network. The macromolecular parameter C.infin. is
incapable of predicting failure strain and toughness in moderately
crosslinked networks while CED can be used with some success in
(meth)acrylate networks. New predictive parameters need to be
developed or previous ones augmented to take into account key
characteristics of network structure. In particular, the
viscoelastic region is of great importance because shape memory
polymers rely on approaching their Tg for actuation. In this
region, both the monomer and network structure play a role in the
large strain properties of the material as was demonstrated
here.
Example 4
The Effect of Chemistry on the Polymerization, Thermo-Mechanical
Properties and Degradation Rate of Poly(.beta.-Amino Ester)
Networks
[0124] PEGDA of four varying molecular weights, M.sub.n258, 302,
575, 700, was used as one diacrylate system. The other diacrylate
system, diol diacrylates (DDA), comprised 1,4-butanediol diacrylate
(DDA198) (Dajac Labs), 1,6-hexanediol diacrylate (DDA226),
1,9-nonanediol diacrylate (DDA268) (TCI). The primary amine was
3MOPA. 2-hydroxy-1-[4-(hydroxyethoxy)phenyl]-2-methyl-1-propanone
(Irgacure 2959) was used as the photoinitiator. All chemicals were
used as received from Sigma Aldrich unless otherwise noted.
[0125] Each of the seven diacrylates was mixed with 3MOPA at molar
ratios from 1.05:1 to 1.25:1 at increments of 0.05. The step-growth
polymerization occurred for 24 hours (h) at 200 rotations per
minute (rpm) at 90.degree. C. on a JKEM reaction block (RBC-20 with
BTS-1500 shaker) to form the macromers. The resulting macromers
were either stored at 4.degree. C. or mixed with 0.5 weight percent
(wt. %) Irgacure 2959 for photopolymerization. The macromers were
placed into a 5 centimeters (cm).times.6 cm.times.0.1 cm teflon
mold, sealed with glass slides, and polymerized for 10 min by a UVP
Blakray lamp (e8 mW/cm2) to form chemically crosslinked networks.
Due to the inherent heterogeneity, the materials were synthesized
and all tested in triplicate, where mean.+-.SD is reported.
[0126] Dynamic mechanical analysis (DMA) of each network was
performed via a TA Instruments DMA Q800. Rectangular samples were
run in tension under strain control of 0.1% according to the
following protocol: equilibrate at -100.degree. C., isotherm for 2
min, ramp 3.degree. C. per minute to 100.degree. C. The glass
transition temperature (Tg) was defined at the peak of the tan d
curve. The molecular weight between crosslinks was calculated
from:
M C = 3 .rho. RT E ##EQU00001##
where E is the modulus, r is the polymer density, R is the gas
constant, T is the temperature in Kelvin, and M.sub.c is the
molecular weight between crosslinks.
[0127] Select networks were analyzed with a TA Instruments DSC
Q100. Samples were cooled at 3.degree. C. per minute to -90.degree.
C., isotherm for 2 min, and then heated at 3.degree. C. per minute
to 100.degree. C.
[0128] AVarianMercury Vx 400.sup.1 HNMR was used to verify the
structures of the macromers after step-growth polymerization in
deuterated chloroform. The spectrum was analyzed via MestRe-C
software to determine the number of hydrogen as well as an estimate
of the diacrylate to amine ratio, and therefore the molecular
weight.
[0129] A Nicolet Nexus 870 FTIR with attenuated total reflectance
(ATR) module was used to characterize the step-growth and
free-radical polymerization. Macromer samples were taken at 2, 4,
8, 16, and 24 hours from the reaction block and the acrylate peak
at 812 cm.sup.-1 was monitored. After 24 hours, the samples were
mixed with 0.5 wt % Irgacure 2959 and polymerized with the UVP
Blakray lamp. The data collection was taken in real time for at
least 5 min to ascertain the degree of conversion to a network from
the peak at 812 cm.sup.-1.
[0130] A sol fraction test was employed to determine the extent of
conversion in network formation. Tert-butyl benzene was used, where
1 square centimeters (cm.sup.2) squares cut from 1 mm thick sheets
were soaked for 48 hours with a change in solvent at 24 hours. The
samples were dried in an oven with dessicant to remove all traces
of solvent and then allowed to equilibrate with the surrounding
atmosphere for 3 days. The sol fraction is defined through
Equation:
SolFraction = 1 - M f M i ##EQU00002##
where M.sub.f is the final mass and M.sub.i is the initial
mass.
[0131] In order to determine the degradation rate and water content
of each material, each material was soaked for varying amounts of
time. Each material was cut from a 1 millimeter (mm) thick sheet
into a 1 cm.sup.2 and placed into a well plate with phosphate
buffered saline (PBS), pH 7.4. The well plates remained in an
incubator at 37.degree. C. on a rotary shaker at 60 rpm. Samples
were patted dry to remove excess water to obtain the wet sample
mass. The samples were dried for 24 h and the mass taken. The water
content of each material is defined by:
WaterContent = M wi M di - 1 ##EQU00003##
where M.sub.wi is the wet mass at time i and M.sub.di is the mass
at time i after 24 hours of drying. The mass loss is defined
by:
Massloss = 1 - M di M o ##EQU00004##
where M.sub.di has been previously described and M.sub.o is the
initial mass.
[0132] Representative DMA and DSC curves are shown in FIG. 31 for
both PEGDA and DDA networks. The glass transition temperature (Tg)
is well below room temperature for all these materials, and thus
the material is rubbery at ambient temperatures. PEGDA700-based
networks have a hump in the modulus above Tg, which signifies
crystallization and subsequent melting due to the high molecular
weight of the PEGDA chain. The crystallization and melting can be
seen in the DSC curves in FIG. 31B for the PEGDA700-based network.
The molecular weight of the PEGDA575-based network or the
DDA226-based network is not high enough to promote crystallization,
thus the lack of the hump in the DMA curves of FIG. 31A and the
lack of crystallization and melting peaks in the DSC curves of FIG.
31B.
[0133] Networks formed from each diacrylate and 3MOPA at varying
molar ratios were tested on the DMA. The modulus in the rubbery
regime at a temperature of Tg+75.degree. C. is given as a function
of the molar ratio as shown in FIGS. 1A and B. In order to compare
rubbery modulus between systems, the rubbery modulus was measured
at the same relative temperature to Tg, Tg+75.degree. C. The
PEGDA302 does not form a network at molar ratio 1.05:1 due to a
lack of acrylate bonds. The modulus ranged from 0.14 to 5.36 MPa
for the PEGDA networks, and from 0.15 to 6.71 MPa for the DDA
networks. The rubbery modulus increases as the molar ratio
increases, as expected from a similar study. However, there is no
obvious trend between the diacrylate molecular weight and the
rubbery modulus. The Tg as a function of the molar ratio is shown
in FIGS. 1C and D. The Tg ranged from -44.3.degree. C. to
-31.degree. C. for the PEGDA networks, and from -50.9 to
-35.6.degree. C. for the DDA networks. The Tg increases as the
molar ratio increases and as the diacrylate molecular weight
decreases for the DDA networks, but the Tg increases only as the
diacrylate molecular weight decreases for the PEGDA networks. The
molecular weight between crosslinks for PEGDA-based and DDA-based
networks for each molar ratio is shown in FIGS. 1E and F. The
molecular weight between crosslinks ranged from 1500 to 115,000
grams per mole (g/mol) for the PEGDA-based networks and from 1200
to 59,000 g/mol for the DDA-based networks.
[0134] A series of structural analyses were performed to help
understand the trends in modulus presented in FIG. 1. In order to
understand the extent of conversion during step-growth
polymerization, the acrylate bond was monitored via FTIR over 24 h.
The molar ratio of 1.20:1 was examined for each diacrylate as shown
in FIG. 32. In FIG. 32A, the PEGDA258 and PEGDA302 converted
quicker and to a higher extent than the PEGDA575 and the PEGDA700
networks. In FIG. 32B, the degree and rate of conversion increased
as the DDA molecular weight decreased. PEGDA and DDA macromers
formed from monomers of similar molecular weight have quite
different step-growth conversions, where PEGDA monomers converted
faster and to a higher degree. Due to the diacrylate to amine ratio
being greater than 1, all of the diacrylate endgroups will not be
completely consumed. When using a molar ratio of 1.20:1 of
diacrylate to amine and all of the amine endgroups react, there
will be a theoretical diacrylate excess of 16.6%; thus, the
expected amount of diacrylate endgroups consumed is 83.3%.
[0135] After 24 h of step-growth polymerization, NMR was used to
verify the chemical structure of the macromers, especially the
presence of acrylate endgroups, and the incorporation of the amine
into the macromer as shown in FIG. 33 by an absence of a peak near
1 parts per million (ppm). From endgroup calculations, an
estimation of macromer molecular weight can be determined. The
average molecular weight of the macromers from NMR for each
backbone chemistry and molar ratio is shown in FIG. 34. The average
molecular weight of the macromers decreased as the molar ratio
increased for both network chemistries. The macromer molecular
weight ranged from 2200 to 37,400 g/mol for the PEGDA-based
macromers, and the macromer molecular weight of the DDA-based
macromers ranged from 1700 to 11,800 g/mol. The macromer molecular
weight of the PEGDA-based macromers showed no direct relationship
with diacrylate molecular weight, where the PEGDA575 and
PEGDA258-based macromers had similar macromer molecular weights and
the PEGDA700 and the PEGDA302-based macromers had similar macromer
molecular weights. The PEGDA258 and PEGDA302-based macromers did
reach a higher conversion during stepgrowth, and therefore have
higher macromer molecular weight. The macromer molecular weight of
DDA-based macromers converged as the molar ratio increased.
[0136] From the FTIR-ATR photopolymerization, the conversion of
remaining acrylate bonds was monitored for networks at a molar
ratio of 1.20:1. The conversion of PEGDA networks and DDA networks
is shown in FIGS. 35A and B, respectively. PEGDA575 and PEGDA700
networks reached high degrees of conversion, while PEGDA258 and
PEGDA302 did not. The DDA226 and DDA268 reached higher degrees of
conversion than the DDA198 network. In essence, the networks formed
from macromers that had high degrees of conversion during
step-growth did not reach high degrees of conversion during
photopolymerization. To compare the conversion measured from
FTIR-ATR photopolymerization, sol fraction testing was conducted
across all molar ratios and diacrylate monomer chemistries and
molecular weights as shown in FIG. 36 PEGDA575 and PEGDA700
networks showed lower sol fractions, 0.03-0.06, respectively,
compared to the PEGDA258 and PEGDA302 networks, 0.13-0.5,
respectively. DDA networks showed similar levels of sol fraction to
the PEGDA networks ranging from 0.04 to 0.38.
[0137] Degradation profiles for all diacrylate molecular weights
and molar ratio of DDA are shown in FIG. 37A. The degradation rates
of the DDA226 and DDA268 are tightly grouped, while the DDA198 is
distinct from the other DDA networks. The marked curves are the
lowest molar ratio, 1.05:1, for each DDA molecular weight. The
water content of the DDA226 and DDA268 does not exceed 1, while the
water content of the DDA198 is higher as shown in FIG. 37B. The
networks with rapid degradation match the networks with high water
content. In the DDA system degradation rate is strongly influenced
by molecular weight but not by ratio for this range of molecular
weights.
[0138] The degradation profiles of PEGDA-based materials are
presented in FIGS. 38A and B. The degradation profiles for PEGDA700
networks for the 5 different molar ratios are shown in FIG. 38A.
The lower PEGDA700 ratios degraded rapidly, while the higher ratios
lasted for at least 24 h. The degradation profiles for all PEGDA
networks at a ratio of 1.10:1 and 1.25:1 are shown in FIG. 38B. All
PEGDA networks at a ratio of 1.10:1 degraded completely within 8 h,
while the networks at a ratio of 1.25:1 lasted 12 h or more. The
water contents for the four PEGDA networks at a ratio of 1.25:1 are
shown in FIG. 38C. The networks with high initial water contents
match the networks with rapid degradation.
[0139] In this study, the effect diacrylate molecular weight,
diacrylate to amine ratio, and diacrylate chemistry on PBAE network
properties was explored. The results show that the degradation rate
and thermo-mechanical properties are greatly influenced. The two
systems, PEGDA-based and DDA-based, were selected for their
different history of use and their diverse chemical properties.
From a biocompatibility standpoint, PEGDA-based polymers are known
for their biocompatibility, and the degradation product containing
3MOPA has yet to be proven harmful. In addition, PEGDA-based
systems were chosen due to previous testing of their mechanical
properties under cyclical loading. As a comparison, DDA-3MOPA
systems were chosen to study the effect of changing the backbone
chemistry from hydrophilic to hydrophobic on the thermo-mechanical
properties and degradation. The molar ratio range was limited to
1.25:1 to prevent non-degradable crosslinks from forming from
excess pure diacrylate. Prior work has demonstrated that the
elastic modulus of the network is affected by changes in macromer
molecular weight by varying the molar ratio. Here we explore the
impact of diacrylate molecular weight and chemistry on step-growth
polymerization, photopolymerization, and subsequent properties.
[0140] Dynamic mechanical analysis was used as a screening method
in order to look at a range of networks of varying crosslinking
density and chemistry. A commonality between all these materials is
their low Tg, as shown by the large drop in storage modulus in
FIGS. 1C and D. The low Tg is due to the lack of steric hindrance
usually created from bulky, rigid side groups, and the enhanced
flexibility from the ethylene glycol, methylene, or amine groups
incorporated into the backbone. Thus, by increasing the diacrylate
molecular weight, more flexible groups are being incorporated
resulting in the subsequent decrease in Tg. In addition, varying
the molar ratio and diacrylate molecular weight and chemistry
produced a broad range of rubbery moduli. The rubbery modulus
increases as molar ratio increases, thus the crosslinking density
increases as the macromer molecular weight decreases due to the
increasing molar ratio as shown in FIGS. 1A and B. However, the
trends in elastic modulus with diacrylate molecular weight were
less obvious and non-monotonic. Low molecular weight diacrylates
may be expected to create denser networks, thus having a higher
modulus. However, in FIG. 1A, the low molecular weight diacrylates
had lower rubbery moduli or in FIG. 1B, the low and high molecular
weight diacrylates had similar rubbery moduli. The crystallization
and melting of PEGDA700-based networks was not expected, but is
possible as the network structure has passed above its Tg in FIG.
31. It also defines a molecular weight boundary, where PEGDA-based
networks do not crystallize when PEGDA is below 700 g/mol. Because
of these unexpected trends in modulus, a further understanding of
both polymerization steps was explored to understand the relation
between diacrylate molecular weight and network modulus.
[0141] The step-growth polymerization of PEGDA-based networks and
DDA-based networks was studied by varying the diacrylate molecular
weight and chemistry while maintaining a constant molar ratio of
1.20:1. The networks synthesized from lower molecular weight
diacrylates, such as PEGDA258 and PEGDA302, converted to a higher
degree and at a faster rate than their higher molecular weight
PEGDA counterparts. The increase of diacrylate molecular weight
decreases the monomer's mobility thus decreasing the rate and
degree of conversion. With a molar ratio of 1.20:1, the conversion
would have been expected to be equivalent for all PEGDA-based
macromers. The higher degree of conversion for the PEGDA258 and
PEGDA302-based macromers may have resulted from termination during
the step-growth polymerization from monomers having only one
acrylate endgroup instead of two or cyclization of the diacrylate
to amine. PEGDA monomers have near 14% impurities comprised of
poly(ethylene glycol) chains and monofunctional poly(ethylene
glycol) acrylate, where over 10% may be monofunctional
poly(ethylene glycol) acrylate. DDA monomers have near 10%
impurities, where 3% comprises the monofunctional diol acrylate.
These monofunctional acrylate impurities will terminate the
step-growth reaction early producing smaller molecules without any
acrylate functionality or cause dangling ends which would not be
elastically effective. The DDA based networks' rates and degrees of
conversion decreased as the diacrylate molecular weight increased
again due to decreased mobility. The effect of decreasing rate and
degree of conversion as diacrylate molecular weight increases is in
agreement with hyperbranched amine-acrylate systems. The macromer
molecular weight post-step-growth polymerization is a key
determinant of the crosslinking density and thus the rubbery
modulus. NMR can provide an estimate based upon the ratio of
acrylate endgroups to amine groups and may count cyclization and
dangling groups with only one acrylate endgroup. By using NMR, the
macromer molecular weight can be compared to the elastically
effective molecular weight between crosslinks via DMA. FIGS. 1F and
33B show similar molecular weights for the DDA-based networks, thus
they are converting ideally and elastically effective chains are
the majority of the chains present. PEGDA575 and PEGDA700-based
networks also show similar molecular weights in FIGS. 1E and 34A.
The PEGDA258 and PEGDA302-based networks have a drastic difference,
where the molecular weight via DMA is much higher. Thus, the
step-growth polymerization is producing elastically ineffective
chains, which would be comprised of a combination of dangling
chains and cycles of low molecular weight as seen in the molecular
weight via NMR. The low molecular weight diacrylates are more
likely to form cycles as seen in kinetic models and other
diacrylate systems, which lowers the formation of crosslinks, thus
lowering their modulus. By comparing the repeating unit structure
of DDA and PEGDA, PEGDA is the more flexible monomer based on its
lower characteristic ratio, where poly(ethylene glycol) and
polyethylene have characteristic ratios of 5.6 and 7.4,
respectively. This increase in flexibility at nearly the same
molecular weight may contribute to the increased cyclization.
Because the step-growth polymerization determines the macromer
molecular weight and the degree of acrylate conversion, this step
will also affect the subsequent polymerization, as will be further
discussed.
[0142] The second polymerization, the UV-photopolymerization, is
responsible for network formation. A critical population of
acrylate endgroups is necessary for full network formation. The
networks that reached high degrees of photopolymerization were the
macromers that did not reach high levels of conversion during their
step-growth polymerization. Thus because the PEGDA575 and
PEGDA700-based networks reached high levels of photopolymerization
conversion, they formed more complete networks and obtained higher
rubbery moduli. By examining the macromer molecular weight in
combination with the degree of conversion, the arrangement of
rubbery moduli in FIG. 1A is made clear. The PEGDA258 and
PEGDA575-based networks have nearly different degrees of
photopolymerization, and thus possess differing rubbery moduli.
Both DDA226 and DDA268-based networks converted to a higher degree
during photopolymerization, but DDA226-based network has a lower
macromer molecular weight than the DDA268, and thus formed networks
with higher rubbery moduli compared with DDA268. The DDA198
converted to a lesser degree during photopolymerization, but still
reached similar values of rubbery moduli as the DDA268. The results
from the sol fraction, an alternative method of measuring network
conversion, are in good agreement with the FTIR-ATR
photopolymerization and the rubbery moduli. The networks suspected
of having dangling endgroups and cycles, PEGDA258 and
PEGDA302-based networks, had the highest sol fractions, thus this
lack of network formation further decreased their modulus
values.
[0143] The results of this study show that the effect of diacrylate
molecular weight and chemistry on the polymerization and mechanical
properties can be fully understood by considering structure after
both the step-growth polymerization and photopolymerization. If the
macromers do not reach high conversion during step-growth but
obtain a high degree of conversion during photopolymerization, then
the effect of diacrylate molecular weight on rubbery moduli can be
understood from the macromer molecular weight. If two macromers
have the same macromer molecular weight via NMR but convert
differently during photopolymerization, then the degree of
conversion during photopolymerization will dictate rubbery modulus.
The stepgrowth polymerization controls the degree of acrylate
conversion necessary for network formation and the macromer
molecular weight that influences crosslinking density. The
photopolymerization controls network formation, but is greatly
influenced by the amount of acrylate endgroups remaining from the
step-growth polymerization.
[0144] The degradation profiles and water content of the networks
are controlled by two different mechanisms. The degradability of
DDA-based networks are affected by their diacrylate molecular
weight as can be seen in FIG. 37A. All five molar ratios of DDA226
and DDA268 have nearly the same degradation profile, and the 5 M
ratios of DDA198 are similar and distinct from the DDA226 and
DDA268 networks. It is clearly seen that as the diacrylate
molecular weight increases, the degradation rate decreases due to a
decrease in water content as shown in FIG. 37B. The water content
follows the same trend as the degradation profiles, where the
number of methylene units or the diacrylate molecular weight is the
controlling factor. The independence of degradation rate from molar
ratio, thus rubbery modulus is unexpected because increasing the
crosslinking density typically alters the degradation rate.
[0145] Unlike the DDA-based networks, the PEGDA-based networks'
degradation profiles are controlled less by their diacrylate
molecular weight, and more by their molar ratio. The low molar
ratios networks are lightly crosslinked thus allowing for large
amounts of water to enter the network, which leads to rapid
degradation. The higher molar ratios of 1.20:1 and 1.25:1 plateau
due to the formation of non-degradable crosslinks and a higher
network density. All PEGDA-based networks follow the same trend,
regardless of the diacrylate molecular weight. All PEGDA-based
networks eventually had water content greater than 500% by the time
of full degradation, which is the main cause for their rapid
degradation. The dual mechanisms illustrate the difference in
backbone chemistry of the diacrylates. Degradation in the DDA-based
networks is more controlled by diacrylate molecular weight while
degradation in the PEGDA-based networks is dominated by molar
ratio. This separation of degradation rate and modulus for the
DDA-based networks is significant, where it will allow for enhanced
tailoring of these networks for tissue scaffolds and drug release
devices.
[0146] While the present disclosure has been described in
connection with a plurality of exemplary aspects, as illustrated in
the various figures and discussed above, it is understood that
other similar aspects can be used or modifications and additions
can be made to the described aspects for performing the same
function of the present disclosure without deviating therefrom. For
example, in various aspects of the disclosure, methods and
compositions were described according to aspects of the presently
disclosed subject matter. However, other equivalent methods or
composition to these described aspects are also contemplated by the
teachings herein. Therefore, the present disclosure should not be
limited to any single aspect, but rather construed in breadth and
scope in accordance with the appended claims
* * * * *